E噩≈S Journal of the European Ceramic Society 22(2002)2777-2787 www.elsevier.com/locate/jeurcerar Intermediate temperature strength degradation in Sic/Sic composites G.N. Morschera.*, J.D. Cawley oHio Aerospace Institute(OAl), NASA Glenn Research Center, MS 106-5, Cleveland, OH 44135, US.A Department of Materials Science and Engineering, Case Western Reserve University, Cleveland, OH 44106, US.A Received 20 November 2001; received in revised form 5 February 2002: accepted 23 February 2002 Abstract Woven silicon carbide fiber-reinforced, silicon carbide matrix composites are leading candidate materials for an advanced jet engine combustor liner application. Although the use temperature in the hot region for this application is expected to exceed 1200C, a potential life-limiting concern for this composite system exists at intermediate temperatures(800=+200C), where sig- nificant time-dependent strength degradation has been observed under stress-rupture loading. A number of factors control the degree of stress-rupture strength degradation. the major factor being the nature of the interphase separating the fiber and the matrix. BN interphases are superior to carbon interphases due to the slower oxidation kinetics of BN. A model for the intermediate temperature stress-rupture of SiC/BN/SiC composites is presented based on the observed mechanistic process that leads to strength degradation for the simple case of through-thickness matrix cracks. The approach taken has much in common with that used by urtin and coworkers, for two different composite systems. The predictions of the model are in good agreement with the rupture data for stress-rupture of both precracked and as-produced composites. Also, three approaches that dramatically improve the intermediate temperature stress-rupture properties are described: Si-doped BN, fiber spreading, and"outside debonding".C 2002 Elsevier Science Ltd. All rights reserved Keywords: BN interfaces: Mechanical properties; Oxidation; SiC/SiC composites 1. Introduction liner I because the"cold side ' of the combustor liner would be exposed to this temperature range, and this Non-oxide ceramic matrix composites(CMCs)such would be the portion of the combustor liner under the as SiC fiber-reinforced SiC matrix composites are envi This is sioned for use as high-temperature,> 1200C, compo- the combustor liner would have to be attached to a nents of gas turbine engines. -However, there exists an metal frame. Therefore, it is important to understand intermediate temperature (600 to 1000 C) regime and predict the time-dependent mechanical behavior at where significant, time-dependent, strength degradation intermediate temperatures for design of these compo- can occur Depending on the application, the inter- sites in components, as well as for finding insights mediate temperature properties may be critical for suc- toward improvement of the intermediate temperature cess of that component. For example, the strength properties. degradation at intermediate temperatures could be an The primary cause for intermediate temperature issue for some applications, e.g. a cooled combustor strength degradation is the oxidation of a non-oxide interphase, usually C or BN, that separates the fibers from the matrix. For C interphases, rapid interphase Corresponding author. Tel: +1-216-433-5512; fax: +1-216-433. removal 5-7 can be associated with fiber strength degra- dation in the form of oxide scale formation 8 or pre- E-mail address: gmorscher(a grc. nasa. gov (G.N. morscher) ferential oxidation of carbon enriched areas on the fiber a By intermediate temperature, we mean an elevated temperature below the supposed use temperature at which the material exhibits a urface. For BN interphases, BN oxidizes to form minima in mechanical properties, i.e. what is sometimes called a liquid B2O3(boria) that leads to the formation of a glass at the interphase region and in the 0955-2219/02/S- see front matter C 2002 Elsevier Science Ltd. All rights reserved. PII:S0955-2219(02)00
Intermediate temperature strength degradation in SiC/SiC composites G.N. Morschera,*, J.D. Cawleyb a Ohio Aerospace Institute (OAI), NASA Glenn Research Center, MS 106-5, Cleveland, OH 44135, USA bDepartment of Materials Science andEngineering, Case Western Reserve University, Cleveland, OH 44106, USA Received 20 November 2001; received in revised form 5 February 2002; accepted 23 February 2002 Abstract Woven silicon carbide fiber-reinforced, silicon carbide matrix composites are leading candidate materials for an advanced jet engine combustor liner application. Although the use temperature in the hot region for this application is expected to exceed 1200 C, a potential life-limiting concern for this composite system exists at intermediate temperatures (800200 C), where significant time-dependent strength degradation has been observed under stress-rupture loading. A number of factors control the degree of stress-rupture strength degradation, the major factor being the nature of the interphase separating the fiber and the matrix. BN interphases are superior to carbon interphases due to the slower oxidation kinetics of BN. A model for the intermediate temperature stress-rupture of SiC/BN/SiC composites is presented based on the observed mechanistic process that leads to strength degradation for the simple case of through-thickness matrix cracks. The approach taken has much in common with that used by Curtin and coworkers, for two different composite systems. The predictions of the model are in good agreement with the rupture data for stress-rupture of both precracked and as-produced composites. Also, three approaches that dramatically improve the intermediate temperature stress-rupture properties are described: Si-doped BN, fiber spreading, and ‘‘outside debonding’’. # 2002 Elsevier Science Ltd. All rights reserved. Keywords: BN interfaces; Mechanical properties; Oxidation; SiC/SiC composites 1. Introduction Non-oxide ceramic matrix composites (CMCs) such as SiC fiber-reinforced SiC matrix composites are envisioned for use as high-temperature, 51200 C, components of gas turbine engines.13 However, there exists an intermediate temperature (600 to 1000 C)a regime where significant, time-dependent, strength degradation can occur.4 Depending on the application, the intermediate temperature properties may be critical for success of that component. For example, the strength degradation at intermediate temperatures could be an issue for some applications, e.g. a cooled combustor liner,1 because the ‘‘cold side’’ of the combustor liner would be exposed to this temperature range, and this would be the portion of the combustor liner under the highest tensile stress. This is especially the case where the combustor liner would have to be attached to a metal frame. Therefore, it is important to understand and predict the time-dependent mechanical behavior at intermediate temperatures for design of these composites in components, as well as for finding insights toward improvement of the intermediate temperature properties. The primary cause for intermediate temperature strength degradation is the oxidation of a non-oxide interphase, usually C or BN, that separates the fibers from the matrix.4 For C interphases, rapid interphase removal 57 can be associated with fiber strength degradation in the form of oxide scale formation 8 or preferential oxidation of carbon enriched areas on the fiber surface.9 For BN interphases, BN oxidizes to form liquid B2O3 (boria) that leads to the formation of a borosilicate glass at the interphase region and in the 0955-2219/02/$ - see front matter # 2002 Elsevier Science Ltd. All rights reserved. PII: S0955-2219(02)00144-9 Journal of the European Ceramic Society 22 (2002) 2777–2787 www.elsevier.com/locate/jeurceramsoc * Corresponding author. Tel.: +1-216-433-5512; fax: +1-216-433- 5544. E-mail address: gmorscher@grc.nasa.gov (G.N. Morscher). a By intermediate temperature, we mean an elevated temperature below the supposed use temperature at which the material exhibits a minima in mechanical properties, i.e. what is sometimes called a ‘‘pest’’ condition.4
2778 G N. Morscher, J. D. Cawley/Journal of the European Ceramic Society 22(2002)2777-2787 matrix crack when the boria reacts with the SiC fibers 2. The process and factors affecting intermediate tem- and matrix 0-lI coupled with B removal from the oxi- perature stress-rupture of woven BN interphase com- dation product as volatile B-containing hydrated species posites in air form in the presence of water vapor. However, the intermediate temperature stress-rupture properties of It was established in Ref. 14 from chemical analysis of Sic/SiC composites with bn interphases have been fiber fracture surfaces(degree of fiber fracture surface shown to be superior to SiC/SiC composites with C oxidation) that whole areas of fibers in a matrix crack interphases when tested in air. failed at the same time during the course of the stress Since BN interphase composites are more durable rupture experiment. It was also discerned from fracture than C interphase composites in oxidizing environments mirror analysis of failed fibers that the amount of at intermediate temperatures, bN has been selected as degradation to the fiber strength was commensurate the interphase material for the earlier mentioned com- with the expected amount from single fiber stress-rup bustor liner application. Therefore, the factors and ture data. 16 That is, no additional strength degradation mechanisms that control intermediate temperature occurred to the population of strongly bonded fibers. stress-rupture for BN interphase composites will be the Therefore, failure under stress-rupture conditions at focus of this study. A model that accounts for some of intermediate temperatures occurs by local overloading these factors will be presented. It will be shown to pre- due to the stress concentration associated with the dict the intermediate temperature rupture data. Finally, strong bonding of fibers to the matrix. Not because the recent enhancements to the microstructure of SiC/BN/ fibers are weakened. The depth of this embrittled region Sic composites resulting in improved intermediate tem- grows as the oxidation front moves deeper into the perature composite performance will be presented. All matrix crack(Fig. 1). Eventually, one of the strongly of the stress-rupture data presented in this work that bonded fibers breaks and causes all the other strongly has not been published earlier was performed in the bonded fibers to fail due to the inability to globally same manner as described in Refs. 14, 15. Woven com- share the increased stress applied to the nearest neigh- posites, 150 mm in length, were tested in tension where bor fibers and unbridged crack growth. This view is the ends of the specimens were"cold-gripped"and a strongly supported by the pattern of fracture mirrors slotted furnace with a 15 mm hot zone was inserted in the from the strongly bonded regions of near fiber contact center region between the grips. The furnace was brought in the micrograph of Fig. l, which are indicative of to temperature prior to the application of the load correlated fiber failure. If the stress transferred to the fiber x(t) BN nnnnnnnnn Embrittled rea (Loc 5)x×a.点k1: O2 and H,0 O and Ho bonded Pristine fiber Area(Global Load Sharing) matrix Fiber break Fig. 1. Idealized schematic representation of oxygen ingress in a matrix crack and an individual fiber failure that leads to failure of all strongly bonded fibers. An example of which is given in the upper left hand corner for a HN/ BN/MI SiC composite
matrix crack when the boria reacts with the SiC fibers and matrix 1011 coupled with B removal from the oxidation product as volatile B-containing hydrated species form in the presence of water vapor.11 However, the intermediate temperature stress-rupture properties of SiC/SiC composites with BN interphases have been shown to be superior to SiC/SiC composites with C interphases when tested in air.9,12,13 Since BN interphase composites are more durable than C interphase composites in oxidizing environments at intermediate temperatures, BN has been selected as the interphase material for the earlier mentioned combustor liner application.1 Therefore, the factors and mechanisms that control intermediate temperature stress-rupture for BN interphase composites will be the focus of this study. A model that accounts for some of these factors will be presented. It will be shown to predict the intermediate temperature rupture data. Finally, recent enhancements to the microstructure of SiC/BN/ SiC composites resulting in improved intermediate temperature composite performance will be presented. All of the stress-rupture data presented in this work that has not been published earlier was performed in the same manner as described in Refs. 14,15. Woven composites, 150 mm in length, were tested in tension where the ends of the specimens were ‘‘cold-gripped’’ and a slotted furnace with a 15 mm hot zone was inserted in the center region between the grips. The furnace was brought to temperature prior to the application of the load. 2. The process and factors affecting intermediate temperature stress-rupture of woven BN interphase composites in air It was established in Ref. 14 from chemical analysis of fiber fracture surfaces (degree of fiber fracture surface oxidation) that whole areas of fibers in a matrix crack failed at the same time during the course of the stressrupture experiment. It was also discerned from fracture mirror analysis of failed fibers that the amount of degradation to the fiber strength was commensurate with the expected amount from single fiber stress-rupture data.16 That is, no additional strength degradation occurred to the population of strongly bonded fibers. Therefore, failure under stress-rupture conditions at intermediate temperatures occurs by local overloading due to the stress concentration associated with the strong bonding of fibers to the matrix. Not because the fibers are weakened. The depth of this embrittled region grows as the oxidation front moves deeper into the matrix crack (Fig. 1). Eventually, one of the strongly bonded fibers breaks and causes all the other strongly bonded fibers to fail due to the inability to globally share the increased stress applied to the nearest neighbor fibers and unbridged crack growth. This view is strongly supported by the pattern of fracture mirrors from the strongly bonded regions of near fiber contact in the micrograph of Fig. 1, which are indicative of correlated fiber failure. If the stress transferred to the Fig. 1. Idealized schematic representation of oxygen ingress in a matrix crack and an individual fiber failure that leads to failure of all strongly bonded fibers. An example of which is given in the upper left hand corner for a HN/BN/MI SiC composite. 2778 G.N. Morscher, J.D. Cawley / Journal of the European Ceramic Society 22 (2002) 2777–2787
G N. Morscher, J. D. Cawley/Journal of the European Ceramic Society 22(2002)2777-2787 pristine weakly bonded fibers cannot be carried by the for a fiber failure originating in the region exposed by a remaining fibers, then the composite fails. Two criteria matrix crack. The number of fibers per tow, number of must be met in order for a composite to fail according tows in a composite cross-section, and the size of the to this process: specimen for a given volume fraction of fibers will affect the total number of fibers in a matrix crack. The effec- 1. A critical number of fibers in a given matrix tive gage length of fibers will be controlled by the ability crack must be strongly bonded to one another or of the fibers to transfer load to the matrix, i. e. interfacial to the matrix. When these fibers fail in a matrix shear strength, the amount of interfacial recession that crack, the stress increase to the remaining may occur, and the number of matrix cracks that are unbroken fibers is sufficient to cause them to fail. exposed in the " hot zone""of the furnace. An increase in 2. An event has to occur to fail one or more of the effective length of fully-loaded fibers will increase those strongly bonded fibers to cause unbridged the likelihood that a strongly bonded fiber will fail in or crack growth. Most likely, this event is caused by near a matrix crack beginning the process of unbridged the failure of one strongly bonded fiber due to crack growth intrinsic fiber strength degradation(flaw growth) of a fiber that is relatively weak in the distribu tion of fiber strengths. It is also possible that 3. A model for intermediate temperature stress rupture fiber-degradation could occur from fiber oxida- of SiC/BN/SiC composites tion depending on the fiber-type and oxidizing environment In order to model this process, an approach con ceptually similar Curtin and coworkers 18-2 approach to model composite strength and individual The kinetics for fiber fusion or the depth into a matrix fiber fracture was employed. Only the simple case of rack away from the exposed surface that fibers are through-thickness cracks was considered. The model strongly bonded depends on the ingress of oxidizing was applied to two SiC fiber Bn interphase MI SiC duction, and the shortest distance between two fibers, for these systems. 4.I g available property information species into the matrix crack, i.e. the rate of oxide pro- matrix systems by usi i.e. the gap that must be filled by oxide. Ingress of oxi- dizing species can only occur if matrix cracks are pre 3. The model sent which intersect load-bearing fibers; therefore, fiber fusion will be dependent on the presence of matrix The stress on the fibers in a bridged matrix cracl cracks, and whether or not those cracks are through the can be found from the applied far-field composite stress, thickness of the specimen. The durability of the inter- 0, and the volume fraction of fibers in the loading phase will affect the rate for fiber-to-fiber fusion. It was direction, f. found for the woven Hi-Nicalon(Nippon Carbon, Co Japan)fiber(HN)reinforced, BN interphase, melt-infil trated (MI) SiC matrix composite of Ref. 14 that the thin carbon layer that exists between the fiber and the bn due to fiber decomposition during matrix processing 7 enhances crack growth and interphase oxidation I5 Also the closer fibers are to one another or the thinner Transfer e interphase coating and the method of interphase used in u= crack coating will be critical. It was found in the earlier study 4 model that over 95% of all the fibers were nearly in contact openIng with one another, i.e. are separated by less than 100 nm R displacement even though the average thickness of the interphase was 0.5 This is nce of woven structures, where the act of weaving tightens tows and forces fibers into intimate contact with one another oeo=o/f Fiber failure depends on the strength-distribution of the fibers in a matrix crack. the number of fibers in a Fig. 2. Schematic representation of stress-profile at and around a matrix crack, and the effective gage length of loaded a composite. sf. m fibers and the matri fibers. A wider distribution of fiber strengths for the rule of The model assumes same average strength will mean a greater probability [Eq(7)
pristine weakly bonded fibers cannot be carried by the remaining fibers, then the composite fails. Two criteria must be met in order for a composite to fail according to this process: 1. A critical number of fibers in a given matrix crack must be strongly bonded to one another or to the matrix. When these fibers fail in a matrix crack, the stress increase to the remaining unbroken fibers is sufficient to cause them to fail. 2. An event has to occur to fail one or more of those strongly bonded fibers to cause unbridged crack growth. Most likely, this event is caused by the failure of one strongly bonded fiber due to intrinsic fiber strength degradation (flaw growth) of a fiber that is relatively weak in the distribution of fiber strengths. It is also possible that fiber-degradation could occur from fiber oxidation depending on the fiber-type and oxidizing environment. The kinetics for fiber fusion or the depth into a matrix crack away from the exposed surface that fibers are strongly bonded depends on the ingress of oxidizing species into the matrix crack, i.e. the rate of oxide production, and the shortest distance between two fibers, i.e. the gap that must be filled by oxide. Ingress of oxidizing species can only occur if matrix cracks are present which intersect load-bearing fibers; therefore, fiber fusion will be dependent on the presence of matrix cracks, and whether or not those cracks are through the thickness of the specimen. The durability of the interphase will affect the rate for fiber-to-fiber fusion. It was found for the woven Hi-Nicalon (Nippon Carbon, Co., Japan) fiber (HN) reinforced, BN interphase, melt-infiltrated (MI) SiC matrix composite of Ref. 14 that the thin carbon layer that exists between the fiber and the BN due to fiber decomposition during matrix processing 17 enhances crack growth and interphase oxidation.15 Also, the closer fibers are to one another or the thinner the interphase, the faster fibers will fuse to one another or to the matrix, respectively. Therefore, the uniformity of the interphase coating and the method of interphase coating will be critical. It was found in the earlier study 14 that over 95% of all the fibers were nearly in contact with one another, i.e. are separated by less than 100 nm, even though the average thickness of the interphase was 0.5 mm. This is especially a consequence of woven structures, where the act of weaving tightens tows and forces fibers into intimate contact with one another. Fiber failure depends on the strength-distribution of the fibers in a matrix crack, the number of fibers in a matrix crack, and the effective gage length of loaded fibers. A wider distribution of fiber strengths for the same average strength will mean a greater probability for a fiber failure originating in the region exposed by a matrix crack. The number of fibers per tow, number of tows in a composite cross-section, and the size of the specimen for a given volume fraction of fibers will affect the total number of fibers in a matrix crack. The effective gage length of fibers will be controlled by the ability of the fibers to transfer load to the matrix, i.e. interfacial shear strength, the amount of interfacial recession that may occur, and the number of matrix cracks that are exposed in the ‘‘hot zone’’ of the furnace. An increase in the effective length of fully-loaded fibers will increase the likelihood that a strongly bonded fiber will fail in or near a matrix crack beginning the process of unbridged crack growth. 3. A model for intermediate temperature stress rupture of SiC/BN/SiC composites In order to model this process, an approach conceptually similar to Curtin and coworkers 1820 approach to model composite strength and individual fiber fracture was employed. Only the simple case of through-thickness cracks was considered. The model was applied to two SiC fiber BN interphase MI SiC matrix systems by using available property information for these systems.14,15 3.1. The model The stress on the fibers in a bridged matrix crack, sf, can be found from the applied far-field composite stress, s, and the volume fraction of fibers in the loading direction, f. Fig. 2. Schematic representation of stress-profile at and around a matrix crack in a composite. sf,m would represent the stress on the fibers where the fibers and the matrix share the load according to the rule of mixtures. The model assumes d/2 extends to sf=0 for simplicity [Eq. (7)]. G.N. Morscher, J.D. Cawley / Journal of the European Ceramic Society 22 (2002) 2777–2787 2779
2780 G N Morscher, J.D. Cawley / Journal of the European Ceramic Society 22(2002)2777-2787 With time, fibers exposed by a matrix crack will fail. Eq (5a) then simply bec Fiber failure is assumed to follow a weibull distribution that can be used to determine the probability for fiber (m+1) failure φ(6)= (8) m+1 P(σ,L)=1-e-中 The fraction of fibers that fail in a matrix crack can be o is the fraction of failed fibers according to determined by summing Eqs.( 8)and(5b)and simplify ing with Eq(4) where m is the Weibull modulus, oo is the reference y, m+/x stress and Lo is the reference length that corresponds to the average fiber strength determined in single fiber tensile tests. L is the effective gage length in the matrix oo(t, T crack. It proved useful to adopt the formulation of urtin's s characteristic stress, Oe, and characteristic where Ks (10) gage length, 8c, where (oc, Sc)=l and Sc is twice the m+1 The operative mechanism to be modeled to R composite rupture is the failure of the strongly fiber that triggers the growth of an unbridged crack fiber slip length, by definition, R is the fiber diameter, through the embrittled fibers in a matrix crack. The time and t is the interfacial shear stress [Eq. (4) dependence for the depth of embrittlement into the The fiber stress around the matrix crack varies composite determines the number of embrittled fibers because of load transfer due to friction(Fig. 2). The available in a matrix crack. The region of fiber embrit- fibers are subject to the maximum fiber stress in the tlement often appears as a"picture frame"4 of strongly crack opening. To determine the total fraction of fiber bonded fibers around the rim of the cross-section of a failures, o can be integrated over the stress transfer specimen fracture surface for a composite with through length, Eo and added to the fraction of fiber failures in thickness matrix cracks. The time dependence for fiber the crack opening width, o to-fiber fusion was determined empirically from the depth of this"picture frame"by examination of rup a(z)"dz (5a) tured-specimen fracture surfaces for two different MI composite systems, one reinforced with HN fibers and =()(∞ Corp, Midland, MI) fibers \4is/mmic (Dow Corning that a semi-empirical parabolic time-dependence was an adequate description: where z is the stress transfer length and u is the crack opening width which can be approximated by: 21 (6) where Cox is an empirical coefficient that best fits the 4Tf2E measured oxidation depth data for a given composite system Since the fiber strength degradation of pu Eq (Sa)was solved by othersfor the case where z is fibers in a through-thickness cracked composite is con- qual to twice the fiber slip length, 8, assuming the far sistent with the measured degradation in fiber strengths field stress on the fibers to be zero. This is an appro- of as-produced fibers, the descriptive expression for time priate assumption because there is a negligible con- dependent fiber strength degradation developed by Yun tribution to from the low far field fiber stress(1/5 and DiCarlo 6 for the latter was used. Their data for ofc).8 Can then be approximated assuming a constant t rupture strength of three Sic type fibers are plotted in from the relationship(see Fig. 2) Fig 4 as a Larson-Miller plot. The conditions for our
f ¼ f ð1Þ With time, fibers exposed by a matrix crack will fail. Fiber failure is assumed to follow a Weibull distribution that can be used to determine the probability for fiber failure: Pð; LÞ ¼ 1e ð2Þ is the fraction of failed fibers according to: ¼ L Lo f o m ð3Þ where m is the Weibull modulus, so is the reference stress and Lo is the reference length that corresponds to the average fiber strength determined in single fiber tensile tests. L is the effective gage length in the matrix crack. It proved useful to adopt the formulation of Curtin’s 18 characteristic stress, sc, and characteristic gage length, dc, where (sc,dc)=1 and dc is twice the c ¼ m o Lo R 1 mþ1 ; c ¼ Rc ð4Þ fiber slip length, by definition, R is the fiber diameter, and t is the interfacial shear stress [Eq. (4)]. The fiber stress around the matrix crack varies because of load transfer due to friction (Fig. 2). The fibers are subject to the maximum fiber stress in the crack opening. To determine the total fraction of fiber failures, can be integrated over the stress transfer length, and added to the fraction of fiber failures in the crack opening width, u: ¼ ðZ dz ¼ ðZ ðzÞ o m dz Lo ð5aÞ u ¼ u Lo f o m ð5bÞ where z is the stress transfer length and u is the crack opening width which can be approximated by:21 u ¼ 2R 4f 2Ef 1 þ Ef f Emð Þ 1f ð6Þ Eq. (5a) was solved by others 22 for the case where z is equal to twice the fiber slip length, d, assuming the far field stress on the fibers to be zero. This is an appropriate assumption because there is a negligible contribution to from the low far field fiber stress (1/5 sfc). Can then be approximated assuming a constant t from the relationship (see Fig. 2): ¼ Rf ð7Þ Eq. (5a) then simply becomes: ðÞ ¼ f c ð Þ mþ1 m þ 1 ð8Þ The fraction of fibers that fail in a matrix crack can be determined by summing Eqs. (8) and (5b) and simplifying with Eq. (4): t;T ¼ f c ð Þ mþ1 m þ 1 þ u Lo f o m ¼ 1 Lo f oðt;TÞ m m þ 1 þ u ¼ K m oðt;TÞ ð9Þ where K ¼ m f Lo m þ 1 þ u ð10Þ The operative mechanism to be modeled to predict composite rupture is the failure of the strongly bonded fiber that triggers the growth of an unbridged crack through the embrittled fibers in a matrix crack. The time dependence for the depth of embrittlement into the composite determines the number of embrittled fibers available in a matrix crack. The region of fiber embrittlement often appears as a ‘‘picture frame’’ 4 of strongly bonded fibers around the rim of the cross-section of a specimen fracture surface for a composite with throughthickness matrix cracks. The time dependence for fiberto-fiber fusion was determined empirically from the depth of this ‘‘picture frame’’ by examination of ruptured-specimen fracture surfaces for two different MI composite systems, one reinforced with HN fibers and the other reinforced with Sylramic (Dow Corning Corp., Midland, MI) fibers 14,15 (Fig. 3). It was found that a semi-empirical parabolic time-dependence was an adequate description: x ¼ Coxt 1=2 ð11Þ where Cox is an empirical coefficient that best fits the measured oxidation depth data for a given composite system. Since the fiber strength degradation of pulled out fibers in a through-thickness cracked composite is consistent with the measured degradation in fiber strengths of as-produced fibers, the descriptive expression for time dependent fiber strength degradation developed by Yun and DiCarlo 16 for the latter was used. Their data for rupture strength of three SiC type fibers are plotted in Fig. 4 as a Larson–Miller plot. The conditions for our 2780 G.N. Morscher, J.D. Cawley / Journal of the European Ceramic Society 22 (2002) 2777–2787
G N. Morscher, J. D. Cawley/Journal of the European Ceramic Society 22(2002)2777-2787 198 One important factor for a fiber strength determina 0.215t tion is the actual strength of the fibers in the composite after processing. The fiber strength of both the Hn and E0.7 SYL fibers at room temperature found by Yun and DiCarlo was 2800 MPa. 6 However, some strength degradation may occur due to composite processing Curtin et al. have established a composite ultimate strength failure criterion based on global load sharing assumptions 2(m+1)=/m+1 for 28> (13) (m+2) Fig 3. Depth of oxidation into the specimen from the surface(face)of the composite versus rupture time for -815C rupture of BN inter- Eq(13)is for matrix crack saturation where pe is the phase, MI SiC composites. The n widths were approximately 2 crack density and pe would be the crack spacing. The m.4, I5 The arrow for the Hi-Nicalon data point was for a specimen room temperature ultimate strengths of all of the com- that did not have through-thickness cracks, i.e. the data point indi- cates that oxygen ingress was at least that deep, three plies, into the posites modeled in this study are known. Therefore, the specimen. 4 The closed symbols indicate 815C experiments where the ultimate strength of the fibers in the composites could specimens failed in the hot zone. The open symbols indicate specimens be estimated by solving for o. from Eqs.(13)and ( that were tested at 960 oC which had failed outside of the hot zone region at a lower temperature estimated to be -870oC 14 m(m+2)R/m+2 study are indicated in Fig 4. This data was best fit, re 2(m+1)Lom+ plotted on a stresss-time plot and re-fitted to fit the common form Eq. 12 is based on the room temperature ultimate (12) fiber strength, Oo(Rn, of 2800 MPa. Assuming the flaw growth mechanism that causes time-dependent fiber where Cr is the coefficient that best fits the fiber rupture strength-degradation rate at intermediate temperatures data and n is the rupture exponent; both are dependent were the same as for single fiber tests and depends on on the fiber type. This then becomes the time-dependent starting flaw size, the time-dependent fiber strength of reference stress for Eq.(9) fibers in the composite can be estimated from Eq . (12) Sylramic 8 Nicalon Hi-Nicalon Yun and Di Carlo [16] 800c1000C1200c 100h100h100h 0.1 500010000150002000025000300003500040000 Larson Miller Parameter, q=T[log t+ 22],(K, hr ig. 4. Rupture strength in a Larson-Miller format for different SiC fibers from Yun and DiCarlo. 6
study are indicated in Fig. 4. This data was best fit, replotted on a stresss-time plot and re-fitted to fit the common form: 0ðt;TÞ ¼ Cf t 1=n ð12Þ where Cf is the coefficient that best fits the fiber rupture data and n is the rupture exponent; both are dependent on the fiber type. This then becomes the time-dependent reference stress for Eq. (9). One important factor for a fiber strength determination is the actual strength of the fibers in the composite after processing. The fiber strength of both the HN and SYL fibers at room temperature found by Yun and DiCarlo was 2800 MPa.16 However, some strength degradation may occur due to composite processing. Curtin et al.19 have established a composite ultimatestrength failure criterion based on global load sharing assumptions: ult ¼ c 2ð Þ m þ 1 m mð Þ þ 2 1 mþ1 m þ 1 m þ 2 ; for 2>1 c ð13Þ Eq. (13) is for matrix crack saturation where c is the crack density and c 1 would be the crack spacing. The room temperature ultimate strengths of all of the composites modeled in this study are known. Therefore, the ultimate strength of the fibers in the composites could be estimated by solving for so from Eqs. (13) and (4): oðcompositeÞ ¼ m mð Þ þ 2 2ð Þ m þ 1 R Lo m þ 2 m þ 1 ult mþ1 " #1 m ð14Þ Eq. 12 is based on the room temperature ultimate fiber strength, so(RT), of 2800 MPa. Assuming the flaw growth mechanism that causes time-dependent fiber strength-degradation rate at intermediate temperatures were the same as for single fiber tests and depends on starting flaw size, the time-dependent fiber strength of fibers in the composite can be estimated from Eq. (12): Fig. 4. Rupture strength in a Larson–Miller format for different SiC fibers from Yun and DiCarlo.16 Fig. 3. Depth of oxidation into the specimen from the surface (face) of the composite versus rupture time for 815 C rupture of BN interphase, MI SiC composites. The specimen widths were approximately 2 mm.14,15 The arrow for the Hi-Nicalon data point was for a specimen that did not have through-thickness cracks, i.e. the data point indicates that oxygen ingress was at least that deep, three plies, into the specimen.14 The closed symbols indicate 815 C experiments where the specimens failed in the hot zone. The open symbols indicate specimens that were tested at 960 C which had failed outside of the hot zone region at a lower temperature estimated to be 870 C.14 G.N. Morscher, J.D. Cawley / Journal of the European Ceramic Society 22 (2002) 2777–2787 2781